Nickel-based alloy

ABSTRACT

The invention provides alloys or superalloys based on nickel essentially comprising the following elements in the amounts indicated as percentages by weight: Cr: 11.5% to 13.5%; Co: 11.5% to 16.0%; Mo: 3.4% to 5.0%; W: 3.0% to 5.0%; Al: 2.2% to 3.2%; Ti: 3.5% to 5.0%; Nb: 0.5% to 2.0%; Hf: 0.25% to 0.35%; Zr: 0 to 0.07%; C: 0.015% to 0.030%; B: 0.01% to 0.02%; and Ni: complement to 100%. The alloy is for use in the production of turbine or compressor disks for turbo-machines, using powder metallurgy techniques.

The invention relates to alloys, or superalloys, based on nickel (Ni)and more particularly intended for the production of compressor orturbine disks for turbo-machines using powder metallurgy processes. Theturbo-machines concerned may be aeronautical (turbojet engine, turbopropengine) or ground-based (gas turbine for the production of energy).

BACKGROUND OF THE INVENTION

In service, compressor and turbine disks, located respectively upstreamand downstream of the combustion chamber of a turbojet engine, aresubjected to mechanical stresses that can be attributed to tension,creep, and fatigue, at temperatures that can reach 800° C. Said disksshould nevertheless have operational service lives of several thousandhours. Thus, said disks must be produced from an alloy which, at hightemperatures, has high resistance to traction forces, very good creepstrength, and crack propagation resistance.

Currently, said disks can be produced from nickel-based alloys usingpowder metallurgy processes, said processes limiting chemicalsegregation phenomena and encouraging good microstructural homogeneityof the alloy.

One known example of a nickel-based alloy is described in Frenchdocument FR-A-2 593 830. Said alloy is sold with reference number N18.

That example of an alloy, along with the alloys of the invention, fallsinto the category of two-phase alloys that comprise: a phase termed thegamma phase formed by a nickel-based solid solution that constitutes thematrix for the metallurgy grains, and a phase termed the gamma-primephase, of structure that is based on the coherent intermetallic compoundNi₃Al. The gamma-prime phase forms several populations of inter- orintra-granular precipitates that appear at different stages of thethermomechanical history of the alloy and that play distinct roles inthe mechanical behavior of the alloy.

It has been shown that the inter-granular precipitate population limitsthe growth of gamma matrix grains during recrystallization heattreatment. Hence, by adjusting the recrystallization heat treatment ofthe alloy, the inter-granular precipitate population, and thus the sizeof said grains, are controlled. Depending on whether the maximumtemperature reached during said heat treatment is higher (supersolvustreatment) or lower (subsolvus treatment) than the solution temperature(or solvus temperature) of the inter-granular precipitates of thegamma-prime phase, recrystallization finishes with a large grain size(for a supersolvus treatment) or a low grain size (for subsolvustreatment).

Tensile strength is generally favored by a reduction in grain size,while creep strength is favored by an increase. Hence, depending on theenvisaged application and the envisaged mechanical characteristics,two-phase alloys are thermomechanically treated to produce either afine-grained microstructure (small grains), i.e. with a grain size ofthe order of 5 μm [micrometer] to 15 μm (i.e. ASTM [American Society forTesting and Materials] indices 12 to 9), or a microstructure with coarsegrains, i.e. with a grain size of the order of 20 μm to 180 μm (i.e.ASTM indices 8 to 2).

Further, the grain strength is ensured by the presence of differentpopulations of intra-granular precipitates of the gamma-prime Ni₃Al basephase and it is generally accepted that the high temperature tensilestrength of said alloys increases with the volume fraction of thegamma-prime phase, said fraction possibly reaching 60%.

The N18 alloy, with a volume fraction of the gamma-prime phase of about55%, principally undergoes subsolvus treatments since a fine-grainedmicrostructure is desirable. The fatigue strength and tensile strengthof said alloy are generally favored over its creep strength, because theservice temperature is often less than 650° C., i.e. relativelymoderate.

At temperatures of more than 650° C., high creep strength is necessaryand, as a result, a coarse-grained microstructure (obtained bysupersolvus treatment) would be better suited. However, carrying out anindustrial scale supersolvus treatment on large diameter disks of N18alloy is very difficult or even impossible because the differencebetween the solvus temperature of the gamma-prime phase and the meltingtemperature (i.e. the onset of melting) of the alloy is too small. Thistemperature range for solution of the gamma-prime phase (i.e. to carryout a supersolvus treatment) is too narrow (less than 30° C.), whichrenders industrial application of the total gamma-prime phase solutionheat treatment uncertain.

Further, high internal stresses arise in the disks during rapid cooling(of the order of 100° C./min) consecutive to total solution heattreatment, and they cause cracks (quench cracks) to appear.

OBJECT AND SUMMARY OF THE INVENTION

The invention aims to provide Ni-based alloys for which it is possibleto carry out not only a subsolvus treatment, but also a supersolvustreatment on an industrial scale and which preferably hashigh-temperature mechanical characteristics, especially creep strength,that are at least equivalent to, and preferably better than those of N18alloy.

This is achieved by alloys that essentially comprise (i.e. apart fromany impurities) the following elements, in the amounts indicated aspercentages by weight:

-   -   chromium (Cr): 11.5% to 13.5%;    -   cobalt (Co): 11.5% to 16.0%;    -   molybdenum (Mo): 3.4% to 5.0%;    -   tungsten (W): 3.0% to 5.0%;    -   aluminum (Al): 2.2% to 3.2%;    -   titanium (Ti): 3.5% to 5.0%;    -   niobium (Nb): 0.5% to 2.0%;    -   hafnium (Hf): 0.25% to 0.35%;    -   zirconium (Zr): 0 to 0.07%;    -   carbon (C): 0.015% to 0.030%;    -   boron (B): 0.01% to 0.02%; and    -   nickel (Ni): complement to 100%.

The Applicant's research that led to the invention shows that theproblems encountered with N18 alloy are linked in part to the highvolume fraction (55%) of the gamma-prime phase in that alloy.

In fact, the Applicant has shown firstly, that said high volume fractiontends to reduce the difference between the solvus temperature of thegamma-prime phase and the melting temperature of the N18 alloy,rendering that difference too small to carry out a supersolvus treatmenton an industrial scale.

Secondly, the Applicant has shown that the internal stresses arising inthe part during rapid cooling consecutive upon total solution heattreatment result in part from precipitation of a high volume fraction ofthe gamma-prime phase.

Finally, the Applicant has shown that when the temperature is held atover 650° C. for a sufficiently long period, the elemental compositionof the N18 alloy causes the development of topologically compact phases,generally denoted sigma and mu phases, which are deleterious to the hightemperature behavior of a disk in operation.

Thus, the composition of the alloys of the invention is selected so asto cause a limited volume fraction of gamma-prime phase to precipitate.

While the alloys of the invention are less rich than N18 alloy ingamma-prime phase, against all expectations, their small-grainedmicrostructure has tensile and creep strength characteristics that arebetter than those of the reference alloy. It also appears that thesealloys have equivalent fatigue-creep crack propagation rates which areequivalent to or even better than those of N18 alloy.

For turbo-machine compressor disks or turbine disks, high tensilestrength is particularly favorable to the rupture behavior of said disksas may occur during accidental overspeeding. This high strength is alsoan indicator of good oligocyclic fatigue properties and adequate servicelives.

Further, the reduction in the volume fraction of the gamma-prime phaserelative to the N18 alloy is favorable to the production of disks havinga coarse-grained microstructure and thus high creep strength at hightemperature (i.e. for temperatures of 700° C. or more). This creepstrength associated with very good tensile and fatigue-creep crackpropagation properties allows these disks to be used at temperaturesthat are higher than in current turbo-machines, providing access tobetter thermal efficiencies and a reduction in the specific consumptionof the turbo-machines.

Production of said coarse-grained microstructure is further facilitatedby the comfortable range of temperatures between the solvus temperatureof the gamma-prime phase and the melting onset temperature for thealloy. Advantageously, the compositions of the alloys of the inventionare such that this range spans 35° C. or more. This means that heattreatments above the solvus temperature can be carried out on anindustrial scale, without risking melting the alloy.

The capability of developing one or the other microstructure,coarse-grained and small-grained, as well as the good mechanicalproperties corresponding to each of said microstructures, is a distinctadvantage in alloys of the invention compared with those in current use,especially N18 alloy.

Further, this capability allows dual-structured disks to be produced. Bycarrying out heat treatment at a temperature gradient, a coarse-grainedstructure is developed in the peripheral zone of the disk where theservice temperatures are the highest and where creep plays a significantrole in material damage, and a small grain structure is developed in thecentral zone of the disk (close to the hub), which is cooler, wheredamage essentially results from traction forces and cyclic stresses.

Despite an aluminum concentration that is lower than that of the N18alloy (which is directly correlated to a smaller volume fraction ofgamma-prime phase), the alloys of the invention have relatively lowdensity, preferably 8.3 kg/dm³ [kilograms/cubic decimeter] or less,which means that the mass of the disk and stresses resulting fromcentrifugal force are limited.

Finally, the elemental compositions of alloys of the invention providethem with good microstructural stability as regards the appearance ofsigma and mu phases, which is retarded to more than 500 hours maintainedat 750° C.

To limit the risk of quench cracking, in particular during treatments ata temperature that is higher than the solvus temperature of thegamma-prime phase, the compositions of the alloys of the invention havea limited gamma-prime phase volume fraction, preferably of 50% or less.Sufficient gamma-prime phase must nevertheless be present, so thegamma-prime phase volume fraction is preferably in the range 40% to 50%.

Advantageously, to obtain said volume fraction of the gamma-prime phasein alloys of the invention, the sum of the Al, Ti, and Nb contents, asatomic percentages, is 10.5% or more, and 13% or less, i.e.10.5%≦Al+Ti+Nb≦13%.

Although precipitation of the gamma-prime phase in Ni-based alloysoccurs exclusively due to the presence of Al in sufficientconcentration, the elements Ti and Nb which, by being substituted forAl, are constituents of that phase, are considered to be elements thatare favorable to the formation of the gamma-prime phase in the sameamount and they are termed gamma-prime-genic. The value of the volumefraction of the gamma-prime phase is thus a function of the sum of theatomic concentrations of Al, Ti, and Nb.

It should be noted that tantalum (Ta) is also a gamma-prime-genicelement, but it does not appear in the composition of the alloys of theinvention. Ta is a high atomic mass element, which means that complexcompositional adjustments have to be made to maintain the density of thealloy within reasonable limits (preferably 8.3 kg/dm³ or less). Further,Ta is expensive and it has not been possible to establish clearly thatit has any beneficial role in crack resistance. Finally, itsstrengthening effect on the gamma-prime phase does not appear to begreater than that of the elements Ti and Nb. It has even been shown thatthe strength of the alloys of the invention is at least equivalent tothat of alloys containing Ta.

Also advantageously, the amounts of Al, Ti, and Nb, as an atomicpercentage in the alloys of the invention, are such that the ratiobetween the sum of the amounts of Ti and Nb and the amount of Al is 0.9or more and 1.1 or less, i.e. 0.9 ([(Ti+Nb)/Al] (1.1.

The Ti and Nb atoms substituting for Al in the gamma-prime phase Ni3Albase strengthen it by mechanisms analogous to those of solid solutionhardening. Said hardening is greater as the ratio [(Ti+Nb)/Al] rises.However, beyond a certain value of the concentration of Ti, the coherentNi3Ti eta phase precipitates in the form of elongate plates that have adeleterious effect on the mechanical behavior, especially on theductility, of alloys containing it. Further, the concentration of Nbmust be limited, since an excessive Nb content is prejudicial to thecrack propagation resistance in this type of alloy.

In accordance with a further aspect of the invention, the amounts of W,Mo, Cr, and Co, as an atomic percentage, are such that the sum of theamounts of W, Mo, Cr, and Co is 30% or more and 34% or less, and suchthat the sum of the amounts of W and Mo is 3% or more and 4.5% or less,i.e.: 30%≦W+Mo+Cr+Co ≦34%; and 3%≦W+Mo ≦4.5%.

The elements which essentially substitute for Ni in the gamma solidsolution are Cr, Co, Mo, and W.

Cr is essential for oxidation and corrosion properties of the alloy, andit participates in hardening the gamma matrix by the solid solutioneffect.

Co improves the high-temperature creep strength of these alloys.Further, an increase in the concentration of Co within the stabilitylimits of the structure of the gamma phase can reduce the solvustemperature of the gamma-prime phase and hence facilitate carrying outthe partial or complete solution heat treatments thereof.

Mo and W greatly harden the gamma matrix by the solid solution effect.However, those elements have high atomic masses and their substitutionfor Ni (in particular substitution of W for Ni) results in a substantialincrease in the density of the alloy.

The amounts of Cr, Mo, Co, and W in the alloys of the invention mustthus be carefully adjusted relative to one another in order to obtainthe desired effects, in particular optimum hardening of the gammamatrix, without in any way risking causing the premature appearance offragile intermetallic compound phases, namely sigma and mu. Said phases,when they develop in excessive quantities, can cause a significantreduction in the ductility and mechanical strength of the alloys.

Finally, it should be noted that the minor elements, which are C, B, andZr, form segregations principally at the grain boundaries, for examplein the form of carbides or borides. They thus contribute to increasingthe strength and ductility of alloys by modifying the chemistry of thegrain boundaries, and their absence would be prejudicial. However, anexcess of those elements causes a reduction in the temperature ofmelting onset and causes excessive precipitation of carbides andborides, which consume the elements of the alloy and which no longerparticipate in hardening the alloy. The concentrations of carbon, boron,and zircon are thus adjusted, in particular with non-zero minimumamounts of carbon and boron, so as to obtain good high-temperaturestrength and optimum ductility for alloys of the invention. Hf is alsopresent in moderate quantities, since that element improves thehigh-temperature inter-granular cracking resistance.

The invention also provides a method of fabricating a part, moreparticularly a turbo-machine part such as a compressor or turbine disk,wherein a blank of said part or the part itself is produced from apowder of an alloy of the invention, using a powder metallurgytechnique.

Advantageously, said blank or said part undergoes recrystallization heattreatment during which the blank or part is brought either to atemperature that is below the solvus temperature of the gamma-primephase of said alloy or to a temperature that is above the solvustemperature of the gamma-prime phase of said alloy, and lower than themelting onset temperature of said alloy, to encourage the development ofa microstructure with a grain size which is adapted to the stressconditions.

BRIEF DESCRIPTION OF THE DRAWING

The invention, its applications and its advantages can be betterunderstood from the following detailed description. Said descriptionmakes reference to the accompanying figures in which:

FIG. 1 is a scanning electron microscope image showing themicrostructure of alloy A, described below; and

FIG. 2 is a scanning electron microscope image showing themicrostructure of alloy C1, described below.

MORE DETAILED DESCRIPTION

The parts produced from the alloys of the invention are preferablyfabricated using powder metallurgy techniques.

As an example, production of a compressor or turbine disk using a powdermetallurgy technique comprises the following steps:

-   -   fabricating a master alloy ingot by mixing and melting metallic        elements that are pure (apart from any impurities);    -   re-melting the ingot and pulverizing it with an inert gas or        remelting the ingot and centrifugal pulverization using a known        rotating electrode technique, to obtain a pre-alloyed powder;    -   consolidating said pre-alloyed powder by hot isostatic pressing        and/or by drawing;    -   forming a disk blank by isothermal forging;    -   heat treating said blank; and    -   Final machining of the disk.

At the end of the isothermal forging, different heat treatment steps maybe selected to obtain the microstructure which is best suited to theenvisaged application. The temperature of the gamma-prime phase solutionheat treatment allows the metallurgy grain size to be controlled:

-   -   with a treatment at a temperature which is below the solvus        temperature of the gamma-prime phase, to obtain a microstructure        with small grains (5 μm to 15 μm); and    -   with a treatment at a temperature in the range between the        solvus temperature of the gamma-prime phase and the melting        onset temperature of the alloy, to obtain a coarse-grained        microstructure (more than 15 μm). Said final treatment can be        carried out industrially only if the difference between the two        said temperatures, termed the “solution window”, is sufficiently        large: for industrial alloys, it is assumed that it must be more        than 30° C., preferably more than 35° C.

The cooling rate which follows the solution treatment can control thedistribution of intra-granular precipitates of gamma-prime phase.

One or more tempering treatments can control the size of the tertiaryprecipitates of gamma-prime phase and relax internal stresses whichresult from quenching.

The nominal compositions of two prior art alloys and three alloys of theinvention, given by way of examples, are shown in Table I in which theamounts of the elements of each alloy are shown as atomic percentages,and in Table II in which the amounts are shown as percentages by weight.Alloys C1, C2 and C3 have a solution window of more than 50° C. and arethus treated using the two types of heat treatment presented above,which provides a great range of microstructures. TABLE I Alloy Co Cr MoW Al Ti Nb Hf C B Zr A 15.0 12.5 3.8 0 9.2 5.3 0 0.125 0.079 0.083 0.022B 12.9 18.1 2.4 1.3 4.6 4.5 0.4 0 0.190 0.077 0.027 C1 15.1 13.6 2.2 1.36.4 5.6 0.5 0.100 0.109 0.093 0 C2 15.4 14.1 2.5 1.5 6.0 5.0 1.0 0.0930.128 0.080 0 C3 12.0 14.6 2.9 1.0 5.5 4.6 1.0 0.100 0.100 0.080 0.038(amounts shown as atomic percentages)

TABLE II Alloy Co Cr Mo W Al Ti Nb Hf C B Zr A 15.9 11.7 6.6 0 4.4 4.5 00.400 0.017 0.016 0.036 B 13.1 16.2 4.0 4.0 2.2 3.7 0.7 0 0.039 0.0140.043 C1 15.4 12.2 3.7 4.0 3.0 4.6 0.8 0.310 0.023 0.018 0 C2 15.5 12.64.1 4.7 2.8 4.1 1.5 0.285 0.026 0.015 0 C3 12.15 13.0 4.8 3.15 2.55 3.81.6 0.310 0.021 0.015 0.060(amounts shown as percentages by weight)

Alloy A is alloy N18 and alloy B is sold with reference numberRené-88DT.

To carry out tests on these alloys, parts were produced by powdermetallurgy using the following procedure:

-   -   fabricating master alloy ingots by mixing and fusing pure        metallic elements;    -   centrifugal spraying with rotating electrodes;    -   consolidating pre-alloyed powders by hot drawing;    -   heat treatment including a subsolvus or supersolvus treatment.

For the subsolvus treatment, a partial solution treatment for thegamma-prime phase was carried out at a temperature below the solvustemperature (Tsolvus) of the gamma-prime phase (at about Tsolvus −25°C.). The rate of cooling was of the order of 100° C./minute aftersolution. This treatment was followed by tempering for 24 hours at 750°C. and air cooling.

For the supersolvus treatment, a total gamma-prime phase solutiontreatment was carried out at a temperature above the gamma-prime solvustemperature (at about Tsolvus +15° C. to 20° C.). The rate of coolingwas of the order of 140° C./min after solution. Said treatment wasfollowed by tempering for 8 hours at 760° C. and air cooling.

Tables III and IV show some results of mechanical tests carried out intension, creep, and crack propagation respectively for alloys whichreceived a subsolvus treatment (Table III) and a supersolvus treatment(Table IV).

The tensile tests were carried out in air at 650° C. for the subsolvustreatment (Table III) and at 700° C. for the supersolvus treatment(Table IV), and Rm corresponds to the maximum stress measured duringthese tests.

The creep tests were carried out in air at 700° C. at an initial stressof 550 MPa (650 MPa [megapascal] for alloy C1). The parameter t_(0.2%)is the time in hours to reach a plastic deformation of 0.2%.

The crack propagation tests were carried out in air and at 650° C. Thestress cycle was as follows: load ramp-up for 10 seconds, hold for 300seconds at maximum load and release in 10 seconds with a load ratio(minimum load/maximum load) of 0.05. The parameter V_(f35) is the crackpropagation rate, measured at a value of delta K of 35 MPa·m^(1/2).TABLE III Tension at 700° C. Creep at 700° C., Crack propagation atAlloy Rm (MPa) 550 MPa t_(0.2%) (h) 650° C. V_(f35) (m/cycle) A 1474 340 12.10⁻⁵ B 1445  610 3.10⁻⁵ C1 1590  3000* 2.10⁻⁵ C2 1635 23003.10⁻⁵ C3 1589 — —*under initial stress of 650 MPa

TABLE IV Tension at 700° C. Creep at 700° C., Crack propagation at AlloyRm (MPa) 550 MPa t_(0.2%) (h) 650° C. V_(f35) (m/cycle) B 1320    1509.10⁻⁶ C1 1440  1750* 3.10⁻⁶ C2 1428 >3000 5.10⁻⁶*under initial stress of 650 MPa

The results of Tables III and IV show that the alloys of the inventioncan produce a large increase in the high-temperature mechanicalproperties (tension and creep) while keeping the crack propagationresistance identical to or better than known alloys.

Referring to FIGS. 1 and 2, micro structural examinations were carriedout on alloys A and C1 which had undergone a subsolvus treatment, todetect the appearance of topologically compact phases (i.e. fragileintermetallic compounds) after an ageing heat treatment of 500 hours at750° C. The observations were carried out by back-diffused electronscanning electron microscopy on non-attacked specimens. In alloy A,severe ageing of 500 hours at 750° C. caused inter- and intra-granularformation of phases rich in heavy elements. These phases show up inclear contrast (white borders) at the grain boundaries in FIG. 1. Thesephases, when formed in excessive quantities, may cause a significantreduction in the ductility and strength of the alloys. Tests on alloy C1which had undergone the same treatment of 500 hours at 750° C. showedthat said phases were not formed during ageing. The alloys of theinvention were thus more stable than alloy A (N18) as regards theformation of fragile intermetallic compounds, which are topologicallycompact phases.

1. An alloy essentially comprising the following elements, in theamounts indicated, as percentages by weight: Cr: 11.5% to 13.5%; Co:11.5% to 16.0%; Mo: 3.4% to 5.0%; W: 3.0% to 5.0%; Al: 2.2% to 3.2%; Ti:3.5% to 5.0%; Nb: 0.5% to 2.0%; Hf: 0.25% to 0.35%; Zr: 0 to 0.07%; C:0.015% to 0.030%; B: 0.01% to 0.02%; and Ni: complement to 100%.
 2. Analloy according to claim 1, wherein the sum of the amounts of Al, Ti,and Nb, as an atomic percentage, is 10.5% or more and 13% or less.
 3. Analloy according to claim 1, wherein the amounts of Al, Ti, and Nb, as anatomic percentage, are such that the ratio between the sum of theamounts of Ti and Nb, and the amount of Al, is 0.9 or more and 1.1 orless.
 4. An alloy according to claim 2, wherein the amounts of Al, Ti,and Nb, as an atomic percentage, are such that the ratio between the sumof the amounts of Ti and Nb, and the amount of Al, is 0.9 or more and1.1 or less.
 5. An alloy according to claim 1, wherein the amounts of W,Mo, Cr, and Co, as atomic percentages, is such that the sum of theamounts of W, Mo, Cr, and Co is 30% or more and 34% or less, and suchthat the sum of the amounts of W and Mo is 3% or more and 4.5% or less.6. An alloy according to claim 2, wherein the amounts of W, Mo, Cr, andCo, as atomic percentages, is such that the sum of the amounts of W, Mo,Cr, and Co is 30% or more and 34% or less, and such that the sum of theamounts of W and Mo is 3% or more and 4.5% or less.
 7. An alloyaccording to claim 3, wherein the amounts of W, Mo, Cr, and Co, asatomic percentages, is such that the sum of the amounts of W, Mo, Cr,and Co is 30% or more and 34% or less, and such that the sum of theamounts of W and Mo is 3% or more and 4.5% or less.
 8. An alloyaccording to claim 4, wherein the amounts of W, Mo, Cr, and Co, asatomic percentages, is such that the sum of the amounts of W, Mo, Cr,and Co is 30% or more and 34% or less, and such that the sum of theamounts of W and Mo is 3% or more and 4.5% or less.
 9. A powder of analloy according to claim
 1. 10. A method of fabricating a part, whereina blank of said part or the part itself is produced from a powder of analloy according to claim 1, using a powder metallurgy technique.
 11. Amethod of fabricating a part according to claim 10, wherein said blankor said part undergoes a recrystallization heat treatment in which theblank or the part is brought to a temperature which is higher than thesolvus temperature of the gamma-prime phase of said alloy and lower thanthe melting onset temperature for said alloy.
 12. A method offabricating a part according to claim 10, wherein said blank or saidpart undergoes a recrystallization heat treatment in which the blank orthe part is brought to a temperature which is lower than the solvustemperature of the gamma-prime phase of said alloy.
 13. A method offabricating a part, wherein a blank of said part or the part itself isproduced from a powder of an alloy according to claim 8, using a powdermetallurgy technique.
 14. A method of fabricating a part according toclaim 13, wherein said blank or said part undergoes a recrystallizationheat treatment in which the blank or the part is brought to atemperature which is higher than the solvus temperature of thegamma-prime phase of said alloy and lower than the melting onsettemperature for said alloy.
 15. A method of fabricating a part accordingto claim 13, wherein said blank or said part undergoes arecrystallization heat treatment in which the blank or the part isbrought to a temperature which is lower than the solvus temperature ofthe gamma-prime phase of said alloy.
 16. A turbo-machine part producedfrom an alloy according to claim
 1. 17. A turbo-machine part accordingto claim 16, having a coarse-grained structure in the zone in which itis subjected to the highest operational temperatures and where creepplays a significant role in damage to the part, and a small-grainedstructure in the zone in which it is subjected to the lowest operationaltemperatures and where damage essentially results from tensile forcesand cyclic stresses.
 18. A turbo-machine part according to claim 17,consisting of a compressor or turbine disk.
 19. A turbo-machine partproduced from an alloy according to claim
 8. 20. A turbo-machine partaccording to claim 19, having a coarse-grained structure in the zone inwhich it is subjected to the highest operational temperatures and wherecreep plays a significant role in damage to the part, and asmall-grained structure in the zone in which it is subjected to thelowest operational temperatures and where damage essentially resultsfrom tensile forces and cyclic stresses.
 21. A turbo-machine partaccording to claim 20, consisting of a compressor or turbine disk.